Process for producing an ODS sintered alloy

ABSTRACT

Process for producing a ductile, high strength, oxide dispersion hardened sintered alloy based on a metal having a high melting point. In the past, oxide dispersion has played only a minor role in comparison with other known processes for increasing strength. The process disclosed permits cost effective production of metallic materials which possess a strength hitherto unattainable by oxide dispersion and a higher ductility than prior art materials. As a result, the metallic and nonmetallic foreign components in the sintered alloy can be restricted to the relatively small quantities of dispersoids and any dissolved residual oxygen. The process consists in an annealing treatment and calls for a specific choice of basis metal and suitable oxide dispersoid. The materials so obtained are used mainly where metallic components possessing high strength and ductility together with a minimal concentration of foreign elements are required, for example in human medicine where stringent requirements concerning corrosion resistance and biocompatibility apply or in nuclear technology to prevent undesirable particle reactions.

BACKGROUND OF THE INVENTION

1. Field of the Invention

The invention concerns a process for manufacturing a ductile, highstrength, oxide dispersion hardened sintered alloy based on a metal witha high melting point, if necessary with small additions of substitutionmixed-crystal phase which, however, do not have a serious effect onalloy properties, in which a metal oxide powder in dispersoid form ismixed with the basic metal powder, using oxides of those metals whosebinding energy at temperatures <0.5 T_(M) is higher than that of theoxides of the basic metal.

2. Description of Related Art

Classical processes for altering the strength properties of metalsinclude the forming of alloys via mixed-crystal phases and mechanicalreshaping. In addition, it is known that the strength of materialsproduced by fusion metallurgy or powder metallurgy can be increased byintroducing or removing dispersoids. According to the definition,dispersoids are particles, usually included in the metallic base matrixin a continuous fashion, which even at higher temperatures do not reactwith the basic metal or dissolve, and are not built into the baselattice as substitution metals. Particularly oxides, carbides andnitrides are used as dispersoids.

According to established doctrine, the disadvantage of dispersionhardening versus alloy hardening by continuous or discontinuousprecipitation of a second phase within the basic phase from a commonsolution (precipitation hardening consists in the fact that "it ishardly possible to achieve the same degree of dispersion and strengthincrease as can be realized with precipitation processes in many cases"(H. Bohm, Introduction to Metallurgy, Hochschultaschenbucher Verlag,Mannheim, Zurich).

For producing dispersion-hardened alloys by processes of powdermetallurgy, the dispersoids are usually introduced by soaking the powderwith a dispersoid suspension, or by blending dispersoids in powder formwith the basic metal powder.

Dispersoids introduced in this manner can be further homogenized by"mechanical alloying". The objective of mechanical alloying is todistribute the dispersoids as homogeneously as possible, even within theindividual metal powder grains. These processes are very time-consumingand require grinding equipment of high quality. They are therefore veryexpensive, and their applicability depends on the state of thecomponents. Moreover, practical application demands a compromise betweenthe degree of homogenization and the cost of grinding, i.e. the grindingoperation is limited in time.

The application DE-A1 35 4 255 contains a proposal for producing an ODSalloy by mixing the basic metal in the form of a salt solution with thedispersion particles in colloidal suspension and to finally reduce it tometal. As a special advantage, the finely distributed, homogeneousintroduction of the dispersoid into the metal matrix is cited. However,even with this process, distribution is limited by the particle size ofthe components.

The production of dispersion hardened alloys consists in introducingparticles as dispersoids which by definition do not react or alloy withthe basic matrix. In connection with this fact, the sintered-metallurgyprocesses for producing dispersion alloys up until the present have useddispersoids with melting points that are usually considerably higherthan the alloy sintering temperature. The dispersoids exist in the solidphase during the entire manufacturing process.

Due to the doctrine mentioned above, that dispersion hardening achievesonly relatively small increases in strength, the additional means ofmixed-crystal alloy hardening or precipitation hardening was applied incases where greater mechanical strength was required. To achieve this,greater doses of additive metals were blended with the basic metals,next to dispersoids.

Next to powder metallurgy processes, it is known that oxide dispersionalloys of high-melting metals can be produced by fusion metallurgy,particularly by arc melting.

For instance, a process is known from DE-C1 12 90 727 for producing aniobium alloy of high strength by adding to the niobium small amounts ofoxygen, carbon and/or nitrogen, plus possibly larger amounts of otherhigh-melting metals, next to 0.5-12% zirconium. This alloy melted in thearc is then solution-annealed at 1600°-2100° for between 5 minutes and 9hours, cooled, reshaped and finally subjected to precipitationannealing. The patent description states that, during solutionannealing, the second phase--meaning the carbides, nitrides and/oroxides contained in the basic matrix (basic metal) after melting--formsa solution with the basic matrix. According to that invention, thesecond phase is to remain in solution during the cold shaping due to thesolution annealing and subsequent quenching, and is to be precipitatedhomogeneously and finely during precipitation annealing. The qualitythat can be achieved is documented by means of examples as well as inthe form of tables of mechanical properties.

According to this patent, the means of substitution mixed-crystalalloying as well as precipitation hardening is used in conjunction withthe means of dispersion hardening, cited in column 1, line 65 of thedescription, for increasing the mechanical strength of such alloys. Thestrength values that are realized are thus the result of two or threestrength- and hardness-increasing processes going on simultaneously.

The relatively small amounts of O, but also N and/or C in the alloyindicate that the precipitation of oxides as a means for increasingstrength plays a relatively minor role in that case. In Example 1,mention is made of remelting the ingot six times in order to assure auseful--but, due to the process used, certainly not good--homogenizationof the metals and dispersoids. Even so, the process is comparativelyexpensive. After melting and also still after hot reshaping, such alloyshave a relatively coarse grain which degrades material strength. Forthis reason, the description in column 1, line 015 etc. expressly warns"not to prolong the solution annealing of the sheet metal unnecessarilyin order to prevent grain growth".

Room temperature data are not given in the description. Experience showsthat in alloys produced with this process, relatively high strength canbe expected, but at the same time ductility at room temperature will below (see e.g. V. G. Grigorovich and E. N. Sheftel, Met. Sci. and HeatTreatment 24 (7-8), p. 472 (1983).

U.S. Pat. No. 3 181 966 describes a basic niobium alloy containing0.25-0.5% oxygen and/or 1-3% zirconium and/or titanium, with a weightratio of oxygen/titanium or zirconium between 3:1 to 12:1. In that case,strengthening of the material is achieved by means of oxide dispersionhardening, plus, corresponding to the examples quoted, also to a certainextent by oxygen in interstitial solution and by alloying niobium withtitanium and/or zirconium. It is pointed out there that higher contentsof oxygen in interstitial solution will cause great brittleness in theniobium. In order to counteract this effect, that process makes use ofmetal oxides of metals having a higher bond energy (negative bondenthalpy) than that of the basic metal only in the presence of excessoxide metal. The additions are added, e.g. as titanium oxide powder andspongelike titanium metal, during an arc remelting process of the highlypurified niobium. The process of cooling, which is important for theform of dispersion precipitation, is paid no attention in the patentdescription. This process does not permit any very fine distribution ofthe dispersoids in the basic metal.

The feasible strength properties of niobium alloys that had undergoneadditional hot reshaping but no recrystallization are summarized in atable and compared with the properties of pure, commercially availableniobium qualities. They will later serve as comparative data for thestrength increases possible with the process of this invention.

The invention at hand has as its objective the development of a processto produce ODS sintered alloys having high ductility and strengthproperties, using a high-melting basic metal, which is more economicalthan known processes. The strength properties of alloys produced withknown metallurgical processes should be at least equaled, both in thedeformed and in the recrystallized state, without making use of theformation of substitution mixed crystals or of the classicalprecipitation of a second metal or compound phase as means to achieveincreased strength.

SUMMARY OF THE INVENTION

The process should make possible very precise control of the extent ofdispersion hardening. The ductility of the alloy should be adequate evenfor subsequent cold shaping of the material.

The properties of a single metallic element, such as its corrosionbehavior and its properties when exposed to radiation, should as much aspossible remain unaffected by foreign elements, and at the same time,the mechanical strength of the metals should be significantly increasedover that of the pure phase, with or without deformation hardening.

According to the invention, this task is accomplished by a process inwhich a pressed blank formed of the powder mixture is sintered, withtemperatures at least temporarily reaching 0.7-0.9 T_(M), while thefollowing processes occur:

the oxide that was introduced is broken down and/or is reduced by thebasic metal, the components which are formed are dissolved in the basicmetal;

the dissolved components are finely distributed in the basic metal dueto diffusion;

part of the total oxygen present in the alloy evaporates in a controlledmanner, preferably as an oxide of the basic metal, from the surface ofthe sintered object.

DETAILED DESCRIPTION OF THE PREFERRED EMBODIMENTS

Due to the characteristics of the invention as stated, the process willbe applicable to but a limited number of alloys. Among the metals havinghigh melting points, primarily those of subgroup V and VI of theperiodic system will be suitable. Due to the free negative bond energy,only a limited number of oxides are usable in each case for the desireddispersion hardening. The table below gives an overview of oxides whichare at least applicable in individual cases and their free bond energy,and for comparison shows the oxides of some high-melting metals havingcomparatively low bond energy values:

                  TABLE                                                           ______________________________________                                                                Approximate value of the                                        Solution metal                                                                              negative free oxide bond                                        oxide and hard,                                                                             energy at 25° C. in                                      temperature-  kilojoules per gram atom                              Solution metal                                                                          resistant metal oxide                                                                       of oxygen                                             ______________________________________                                        Silicium  SiO.sub.2     403                                                   Titanium  TiO.sub.2     424                                                   Zirconium ZrO.sub.2     512                                                   Aluminum  Al.sub.2 O.sub.3                                                                            529                                                   Beryllium BeO           584                                                   Thorium   ThO.sub.2     613                                                   Chromium  Cr.sub.2 O.sub.3                                                                            348                                                   Magnesium MgO           572                                                   Manganese MnO           365                                                             MnO.sub.2     233                                                   Lanthanum La.sub.2 O.sub.3                                                                            580                                                   Hafnium   HfO.sub.2     566                                                   Barium    BaO           529                                                   Strontium SrO           560                                                   Calcium   CaO           605                                                   Yttrium   Y.sub.2 O.sub.3                                                                             604                                                   Niobium   Nb.sub.2 O.sub.5                                                                            357                                                   Tantalum  Ta.sub.2 O.sub.5                                                                            388                                                   Vanadium  VO            416                                                   Molybdenum                                                                              MoO.sub.2     251                                                             MoO.sub.3     227                                                   Tungsten  WO.sub.2      251                                                             WO.sub.3      247                                                   Rhenium   ReO.sub.3     189                                                   ______________________________________                                    

An essential factor which governs the choice of suitable combinations ofbasic metal and dispersoid from case to case is the solubility of theoxygen and the oxide metal in the basic metal at the applicablesintering temperature, as well as the melting point of the oxide metalitself. A solubility which is too low, or the formation of intermetalliccompounds between the oxide metal and the basic metal, preclude certaincombinations of metal and oxide, or at least limit the attainabledispersoid percentage in the alloy.

All of the three processes which proceed concurrently during solutionannealing according to the invention are known in themselves, and so arethe means and measures to assure a well-controlled implementation ofthese processes. It is therefore feasible for the average professionalto take appropriate steps in each specific case to attain the desiredbalancing between the three processes.

The concentration of the oxide in the basic metal essentially determinesthe temperature at which the various processes specified by theinvention occur, or become dominant compared to the others. By adaptingthe sintering time and temperature to the components present in thealloy at hand, as well as to their concentration, it is possible toachieve the three processes for oxide homogenization in the course ofthe sintering process.

The total oxide content in the sintered material should preferably beset in such a way that only the exact stoichiometric amount required forforming the oxide remains, which in the strict sense is valid only forthe center of the sintered object due to a diffusion-controlledconcentration profile. In certain cases, the oxide content will be setto a lower value, i.e. below the stoichiometric level, in order toprevent an excessively rapid--and thus usuallycoarse-grained--precipitation of the oxide during cooling after theannealing treatment; this at the expense of a slight reduction instrength.

An excess of oxygen in the material to be sintered will lead to oxygenin interstitial solution next to completely precipitated oxide. Anoxygen deficit results in incomplete oxide precipitation. In the lattercase, part of the oxide metal remains in solution in the basic matrixand will therefore act as a getter for impurities--but also as amixed-crystal component.

Since any excess oxygen in interstitial solution which is notprecipitated as oxide will result in an added increase in strength onthe one hand, but on the other also causes a decrease in ductility,practical application demands that an optimum of all influencing factorsbe determined which takes all requirements into account.

The sintering and annealing process can be carried out by means ofdirect sintering as well as by indirect sintering. In the directsintering process, the material to be sintered is heated by a directpassage of current. The required water cooling of the connectors permitsan especially rapid cooling of the material to be sintered when thesintering process is ended.

Subsequent to the sintering process with solution annealing, theprecipitation in the form of very fine, homogeneously distributed oxideparticles will already occur during the cooling phase, or during asubsequent precipitation annealing step, depending on the dispersoid andits concentration. In this process, the rapidity of cooling plays animportant part, the more so the higher the oxide concentration in thealloy. Directly sintered material can be quenched to low temperaturesparticularly quickly. By heating the alloy, e.g. before extruding as afirst reshaping process, the precipitation of the oxide particles is incertain cases made possible in the first place, or is made complete.

In order to apply mechanical reshaping processes, especially coldshaping by forging, rolling or hammering, the oxide dispersion alloyaccording to the invention must have adequate ductility in addition tohigh strength. It is therefore important to position the strengthproperties of the alloy according to the invention as closely aspossible to a limit which can still just be tolerated, by choosing thedispersoid concentration, but above all by correct control of thesolution annealing step according to the invention.

According to a preferred form of implementation of the invention, thealloy consists of niobium or tantalum as a basic metal and contains,next to small amounts of oxygen in solution, essentially 0.2-1.5% byweight of oxide, using one or more of the metals Ti, Zr, Hr, Ba, Sr, Ca,Y, La.

Particularly outstanding results are obtained with a niobium alloycontaining 0.2-1% by weight of titanium and oxygen, where, next to smallamounts of oxygen in interstitial solution in the niobium basic matrix,TiO₂ is present as a finely distributed dispersoid in the basic matrix.Another preferred niobium alloy contains 0.2-1.5% by weight of ZrO₂.

It was surprising, and not predictable in the extent it occurred, tofind the unusually high strength values achieved through the invention,at a comparably high ductility for dispersion sintered alloys. Forinstance, in the publication "Niobium, TMS-AIME, Proceedings of theInternational Symposium 1981, ed. H. Stuart (1984)" it is stated on page247 that dispersion hardening in Niobium can be attained only to a veryslight extent due to the lack of dispersoids with a sufficiently finedistribution. Even in those cases where alloys were produced usingfusion metallurgy with annealing methods roughly comparable as totemperature and time, the results of the invention at hand could noteven be approximated. Rather, it has to be assumed that due to thedifferent conditions prevailing, in the other process the threeprocesses which occur next to sintering in the invention at hand cannotbe balanced with each other in a comparable manner. In particular, themetal component of the disintegrating oxide in melted alloy materialwill evaporate much more easily from the alloy, compared to sinteredalloys. It therefore does not get a chance to disperse in the basicmetal relatively homogeneously.

To the extent that oxide dispersion alloys have so far been produced bymeans of sintering, the sintering process took place at relatively muchlower temperatures than in the invention at hand. This was to make surethat the oxide particles distributed within the stamped part wouldremain in the place where they were introduced with as little change aspossible and stationary.

It was surprising that the annealing treatment according to theinvention was indeed feasible to the extent actually achieved. Accordingto prevalent doctrine, it had to be feared that, at the annealing andsintering temperatures utilized in the invention, the dissolved oxidemetals would also evaporate from the surface of the sintered object athigh rates, next to the oxides of the basic metal. For, if the requiredconditions for oxide bond energies are met, the melting points of theoxide metals can be significantly lower that the desirable annealingtemperatures according to the invention, and are indeed lower than theannealing temperatures in preferred forms of implementation.

A significant advantage of the process according to the invention is itseconomy. To the extent that dispersion alloys have until now beenproduced by fusion metallurgy, including roughly comparable annealingmethods, the total manufacturing process has been significantly morecost-intensive--e.g. due to melting and remelting of the oxides in theingot by means of arc melting--while the strength gain was clearly less.

Based on the strength values that can be achieved by the variousprocesses, it may be assumed that ODS sintered alloys according to theinvention will achieve much finer oxide particles and homogeneousdispersoid distributions in the basic matrix than with conventionalfusion metallurgy processes including an annealing treatment. As anadditional advantage, sintering consistently yields a much finer grainthan fusion metallurgy.

A significant economic advantage of the process according to theinvention stems from the integration of the annealing treatmentaccording to the invention into the overall sintering process required.

Comparable high-strength alloys--especially at room temperature andmedium elevated temperatures--have until now been obtained for the basicmetal in question only by the formation of a mixed-crystal phase, insome cases with the precipitation of a second metal phase. Intentionallyomitting the formation of a mixed-crystal phase has the followingadvantages:

ODS sintered alloys have comparably high ductility and can therefore bereshaped much more economically to achieve higher final strength values;

these alloys are consistently more corrosion-resistant than thoseproduced by known processes;

typical properties of individual basic metals which are essential fortheir applicability, such as extreme corrosion resistance and thereforebiocompatibility in the case of human implants, but also the use of e.g.niobium due to its low neutron capture cross-section, are practicallyunaffected by the low dispersoid concentrations.

Materials produced according to the process described are required inchemical manufacturing just as much as in tools for high-speed shapingof special alloys, such as super alloys.

An important field of application of niobium and tantalum alloys existsin implants in human medicine. The use of such extremely pure niobiumand tantalum alloys, which are known to be especially compatible withhuman tissue, until now has failed in many cases due to theirinsufficient strength properties. Niobium and tantalum alloys producedby the process according to the invention therefore broaden theapplication area in implant medicine considerably.

A promising application area for alloys according to the invention liesin piping systems for alkaline metal cooling circuits, such as innuclear plants.

The excellent strength properties of alloys according to the inventionwill be illustrated in conjunction with the following examples.

EXAMPLE 1

An alloy of niobium with 0.5% TiO₂ by weight is produced by the processaccording to the invention. For this purpose, 3980 grams of niobiumpowder having a mean grain size of 10 μm and an oxygen content of <1000ppm is blended homogeneously for one hour with 20 grams of TiO₂ powderagglomerate having a mean grain size of 0.25 μm.

This powder mixture is then pressed hydrostatically at about 2000 bardown to 80% of theoretical density.

The pressed object thus obtained is heated slowly under a high vacuum(less than 1×10⁻⁵ mbar) and is finally sintered for 12 hours at atemperature of 2100° C. These sintering conditions are geared to thesize of the samples and to the diffusion and degassing processes to berealized. This leads to a disintegration and the formation of a solidsolution of the TiO₂, as well as to the diffusion of the Ti and O₂components in the niobium. In addition, part of the oxygen is evaporatedfrom the surface of the sintered object, primarily in the form ofniobium oxide.

This results in a very homogeneous distribution of titanium and oxygen,achieving a stoichiometric proportion in the core area of the sample,and in a slightly sub-stoichiometric proportion as to oxygen in theperipheral area of the sample. Further, it was found that theconcentration of the titanium over the entire cross-section of thesintered object is nearly constant, except for a border zone in the mmarea.

Due to the low concentration of TiO₂ in the alloy, there is nosignificant precipitation of TiO₂ during the cooling period followingthe sintering step, but a nearly complete precipitation during apreheating and precipitation step of about one hour at the start of thehot reshaping process. Electron microscopic analyses of samplesfollowing the precipitation annealing step showed that the alloycontained very homogeneously distributed, fine-grained TiO₂ particleswith a particle size of 2-20 nm, predominantly in the range 8-12 nm.

Such alloys can be further processed by the known hot and cold reshapingprocesses. In the case at hand, the first step is a hot reshaping byextrusion at 1000° C. with a reshaping ratio of 8.7:1. The alloy samplewas then processed further by profile rolling and round hammering to acold reshaping factor of 72%. It was possible without any problem toincrease the cold reshaping factor up to 99.9% without intermediateannealing.

Strength tests were then conducted with standard samples made of rods of8 mm diameter. The resulting strength values are summarized in Table 1under Position 1. The table shows two sets of tensile strength at roomtemperature, 800° C., 1000° C. and 1200° C., both for the deformedsample and after recrystallization for 1 hour at 1400° C. The tablecontains the appropriate elongation values next to the tensile strength.

Next to tensile strength, the fatigue strength of such alloys was alsotested. The measurements, using an ultrasonic method, showedabove-average results, with a fatigue strength of about 400 N/mm² in airafter 2×10⁸ cycles.

The alloy possesses excellent ductility. This shows up, for one thing,in excellent machinability, and also in a very low transitiontemperature of about -50° C., a high notch impact strength of about 135J/cm² at room temperature and a high breaking elongation of >10% withdeformed material.

EXAMPLE 2

An oxide dispersion hardened niobium-1 TiO₂ alloy was produced, usingthe process described in Example 1. Twice as much TiO₂ was added as inExample 1.

In contrast with Example 1, a partial precipitation of TiO₂ was observedin this case even during the cooling period following the sintering andreaction annealing process. When the alloy was preheated prior to thehot reshaping process, the titanium still in solution was precipitatedpractically entirely as TiO₂.

The increased TiO₂ content of the alloy caused a higher deformationresistance, so that the samples can better be annealed in between theindividual steps of cold reshaping in order to attain a more evenstructure.

The tensile strength and elongation measured with this sample are shownin Table 1 under Position 2.

EXAMPLE 3

A Niobium-0.5 ZrO₂ alloy was produced according to the process stepsdescribed in Example 1.

Particularly in view of the rapid cooling of the sintered material aftersintering and annealing, the pressed powder blank was processed furtherby way of direct sintering.

Since ZrO₂ is more stable than TiO₂, the sintering temperature wasincreased to 2300° C. in order to assure on the one hand that the ZrO₂components would dissolve completely, but on the other hand also toobtain a somewhat lower total oxygen content of the sample so as toprevent an overly rapid and comparatively coarse re-precipitation of theoxide during the cooling of the sample following the sintering process.A rapid cooling of the sintered object was assured by known measures.

Taking into account the higher stability of ZrO₂ as compared to TiO₂,the preheating or precipitation temperature preceding the first hotreshaping step was increased by 100° C. to 1100° C.

Further process steps were carried out corresponding to Example 1.

The tensile strength and elongation data in the reshaped as well as inthe recrystallized state are shown in Table 1 under Position 4.

Table 1 shows in Positions 1 through 7 the tensile strength andassociated elongation data at various temperatures for a number ofdifferent samples.

The samples are:

Position 1 an Nb-TiO₂ alloy as described in Example 1 of the inventionat hand

Position 2 an Nb-TiO₂ alloy as described in Example 2 of the inventionat hand

Position 3 an Nb-1.5 Ti-0.5 O alloy as described in U.S. Pat. No.3,181,945, quoted with respect of the state of the art

Position 4 an Nb-ZrO₂ alloy as described in Example 3 of the inventionat hand

Position 5 an Nb-1 Zr alloy according to the state of the art ("Niobium,TMS-AIME Proceedings of the International Symposium 1981")

Position 6 an Nb-1 Zr-0.25 O alloy as described in U.S. Pat. No.3,181,945, quoted with respect of the state of the art

Position 7 a very pure Niobium material according to values cited in theliterature and our own measurements.

The results according to the invention cannot be entirely compared withthe values cited in the literature, since, for one thing, thedeformation process of the samples according to the state of the art asquoted was not described in detail, and also because on the basis of thedetailed description given, it must be assumed that next to oxidedispersion precipitates the alloy also still contained significantpercentages of oxide metals of the dispersion oxides in the basicmatrix, exerting an alloy effect which acts to increase strength.

However, it can be stated purely in a qualitative sense thatstate-of-the-art technology cannot produce strength values comparable tothose of the invention at hand. The data for pure Niobium in Position 7show that dispersion alloys produced according to this invention canattain much higher strength properties, at least at room temperature,than by reshaping and possibly recrystallizing pure Niobium.

                  TABLE 1                                                         ______________________________________                                                                    Test  Tensile                                                                              Elon-                                     Material               temp. str.   gation                               Pos. % by weight   State    °C.                                                                          MP a   %                                    ______________________________________                                        1    Nb--0.5 TiO.sub.2                                                                           reshaped RT    950    12                                                                800  405    12                                                               1000  350    12                                                               1200  250    15                                                      recryst. RT    490    34                                                                800  175    33                                                               1000  135    46                                   2    Nb--1 TiO.sub.2                                                                             reshaped RT    1100   12                                                      recryst. RT    535    29                                   3    Nb--1.5 Ti--0.5 O                                                                           hot      RT    506    29                                        (US-PS 3 181 946)                                                                           reshaped  871  307    19                                                                982  251    14                                                               1204  185    20                                   4    Nb--0.5 ZrO.sub.2                                                                           reshaped RT    760    11                                                      recryst. RT    450    32                                   5    Nb--1 Zr      reshaped RT    350-550                                                                              5-15                                      (Niobium      recryst. RT    290    35                                        TMS-AIME)               800  190    18                                                               1000  135    32                                                               1200   90    77                                   6    Nb--1 Zr--0.25 O                                                                            hot      RT    530    16                                        (US-PS 3 181 946)                                                                           reshaped  982  312    17                                                               1093  224    26                                   7    Niobium pure  reshaped RT    300-550                                                                              2-15                                                    recryst. RT    200-300                                                                              20-45                                8    Ta--0.5 TiO.sub.2                                                                           reshaped RT    890    13                                                      recryst. RT    470    31                                   9    Tantalum pure reshaped RT    450-650                                                                              2-7                                                     recryst  RT    300-350                                                                              35-55                                ______________________________________                                    

EXAMPLE 4

Analogous with implementation Examples 1 through 3, an alloy is producedconsisting of tantalum and 0.5% by weight of TiO₂, where the highermelting point of tantalum has to be taken into account for some of theprocess parameters.

7760 grams of tantalum powder with a mean grain size of 9.5 μm, havingan oxygen content of 1050 ppm, are homogeneously blended with 39 gramsof TiO₂ with a mean grain size of 0.25 μm (the identical oxide powder asin Examples 1 through 3).

In order to avoid too much of an O₂ loss due to evaporation of tantalumsuboxides (TaO, TaO₂), the sintering temperature is set to 2300° C.instead of the usual ca. 2600° C. In this manner, a nearlystoichiometric oxygen concentration is attained, corresponding to thetitanium concentration as introduced. The lower sintering density due tothe lower sintering temperature is entirely sufficient for completepacking during the subsequent extrusion step. The precipitationannealing step for precipitating very fine TiO₂ particles is preferablycarried out at 1100° C. in this case.

Due to the high hot strength of the tantalum, extrusion is done at 1200°C. The cold reshaping which follows is carried out by means of profilerollers and round hammering for a total reshaping factor of about 80%.

Under Position 8, Table 1 shows the tensile strength and elongationvalues in the reshaped state and after recrystallization, again obtainedwith 8 mm test rods. The high recrystallization temperature (1600° C., 1hour) leads to a plainly visible coarsening of the TiO₂ dispersoids andthus to a weakening of the dispersion hardening compared to thecold-reshaped material. The combination of cold reshaping and dispersionhardening thus results in especially high strength values whileretaining adequate ductility. For comparison, Position 9 shows thevalues for pure tantalum at 82% reshaping, while the manufacturing stepsand process parameters correspond to those named above.

We claim:
 1. A process for producing a ductile, high-strength, oxidedispersion hardened sintered (ODS) alloy of a base metal having a highmelting point (Tm), comprising:forming a powder mixture by blending apowdered form of said base metal with a dispersoid comprised of a metaloxide powder, said metal oxide powder possessing a higher bond energyvalue than the oxides of said base metal at temperatures less than 0.5Tm; pressing said powder mixture into a pressed blank form; andsintering said pressed blank form at temperatures reaching 0.7-0.9 Tmsuch that said dispersoid is decomposed into its constituent components,and said constituent components are homogeneously dispersed throughoutsaid base metal.
 2. A process for producing a ductile, high-strength,oxide dispersion hardened sintered (ODS) alloy according to claim 1,wherein said ODS alloy includes a small percentage of a mixed-crystalphase of said base metal.
 3. The process according to claim 1, whereinsaid sintering step includes evaporating part of the oxygen present insaid ODS alloy from the surface of said pressed blank as an oxide ofsaid base metal.
 4. An ODS alloy produced according to claim 1, whereinsaid base metal comprises a metal from the group consisting of niobiumor tantalum;said ODS alloy contains a small percentage of oxygen; andsaid metal oxide powder consists of in the range of 0.2-0.5% by weightof said ODS alloy of an oxide of a metal from the group consisting ofTi, Hf, Ba, Sr, Zr, Ca, Y, or La.
 5. An ODS alloy produced according toclaim 1, wherein said base metal comprises niobium;said alloy containsdissolved oxygen; and said metal oxide powder consists of TiO₂ in therange of 0.2-1% by weight of said ODS alloy.
 6. An ODS alloy producedaccording to claim 1, wherein said base metal comprises niobium;saidalloy contains dissolved oxygen; and said metal oxide powder consistingof ZrO₂ in the range of 0.2-1.5% by weight of said ODS alloy.
 7. An ODSalloy produced according to claim 3, wherein said base metal consists ofmetal from the group consisting of niobium or tantalum;said ODS alloycontains a small percentage of oxygen; and said metal oxide powerconsists of in the range of 0.2%-0.5% by weight of said ODS alloy of anoxide of a metal from the group consisting of Ti, Hf, Ba, Sr, Zr, Ca, Y,or La.
 8. An ODS alloy produced according to claim 3, wherein said basemetal comprises niobium;said alloy contains dissolved oxygen; and saidmetal oxide powder consists of TiO₂ in the range of 0.2-1% by weight ofsaid alloy.
 9. An ODS alloy produced according to claim 3, wherein saidbase metal comprises niobium;said alloy contains dissolved oxygen; andsaid metal oxide powder consisting of ZrO₂ in the range of 0.2-1.5% byweight of said ODS alloy.